Semiconductor devices, all types of which are related to the present invention, are used in a wide variety of electronic applications. Semiconductor devices include diodes, transistors, integrated circuits, sensors, and opto-electronic devices, such as light-emitting diodes, diode lasers, radiation detectors, and solid-state photodetectors. Various semiconductor devices using silicon or compound semiconductors, such as gallium arsenide (GaAs) and gallium phosphide (GaP), are commonly used. In order to fabricate semiconductor devices with desired operational electrical performance characteristics, it is necessary to be able to grow high-quality, low-defect-density single crystal films with controlled impurity incorporation while possessing good surface morphology. The substrate upon which the film is grown should also be a high-quality, low-defect-density single crystal as crystal film quality is well-known to depend upon substrate quality to a large extent. In recent years, there has been an increasing interest in research on wide-bandgap (WBG) semiconductors for use in high temperature, high power, high frequency, and/or high radiation operating conditions under which silicon and conventional Group III-V semiconductors cannot adequately function. Particular research emphasis has been placed on SiC, and Group III-Nitride compounds and their alloys, including AlN, GaN, InGaN, AlGaN, and others. Diamond is also a potentially important WBG material.
Conventional semiconductors are unable to meet some of the increasing demands of the automobile and aerospace industries as they move to smarter and more electronic systems. New wide-bandgap (WBG) single-crystal semiconductor materials are being developed to meet the diverse demands for more power at higher operating temperatures. Two of the most promising emerging wide bandgap semiconductors are silicon carbide (SiC) and gallium nitride (GaN). The bandgap of these materials is two to three times as large as that of silicon. This advantage theoretically translates into very large improvements in power handling capabilities and higher operating temperatures that will enable revolutionary product improvements. Once material-related technology obstacles are overcome, SiC's properties are expected to dominate high power switching and harsh-environment electronics for manufacturing and engine control applications, while GaN will enable high-power, high-frequency microwave systems at frequencies beyond 10 GHz. To date, the best SiC devices to our knowledge are fabricated in homoepitaxial films while Group III-Nitride (also referred to herein as III-Nitride) devices are mostly fabricated in heteroepitaxial films (i.e., grown on SiC or sapphire wafers with electrically-active device layers of GaN, AlGaN, AlN, etc.) because production of bulk III-Nitride wafers is currently in an early stage of development.
The performance and the commercialization of high-electric-field SiC power devices is well known to be severely limited by the presence of dislocations in the SiC substrate that, until now, have propagated into the epitaxial layers making up the SiC devices. The best performing SiC high field devices have always been those that are small enough to fit between crystal dislocations, since device performance degrades as the sizes of the devices increase to encompass more and more dislocations. This performance degradation is more fully described in Chapter 6, of the VLSI Handbook, edited by Wai-Kai Chen, CRC Press LCC, Boco Raton, Fla., 2000. One kind of dislocation found in abundance in SiC substrates and epilayers are screw dislocations. Screw dislocations in SiC are further described in three technical articles of M. Dudley et al, “Quantitative Analysis of Screw Dislocation in 6H—SiC Single Crystals,” II Nuovo Cimento, Vol. 19D, No. 2-4, pp 153-164, 1997 (Presented at the 3rd European Symposium on X-ray Topography and High-Resolution X-ray Diffraction (X-TOP '96, 22-24 Apr. 1996, Palermo, Italy); J. A. Powell, P. Pirouz, and W. J. Choyke, “Growth and Characterization of Silicon Carbide Polytypes for Electronic Applications,” Chapter 11 in Semiconductor Interfaces, Microstructures, and Devices: Properties and Applications, Z. C. Feng, Ed. Bristol, United Kingdom: Institute of Physics Publishing, 1993, pp. 257-293, and P. G. Neudeck and J. A. Powell, “Homoepitaxial and Heteroepitaxial Growth on Step-Free SiC Mesas,” in Silicon Carbide-Recent Major Advances, W. J. Choyke, H. Matsunami, and G. Pensl, Eds. Heidelberg, Germany: Springer-Verlag, 2003, pp. 179-205. Additionally, processes for relocating screw dislocations are described in the already mentioned technical article of P. G. Neudeck et al and U.S. Pat. No. 6,783,592B2.
Silicon carbide crystals exist in hexagonal, rhombohedral and cubic crystal structures called polytypes. A general discussion of SiC crystal structure can also be found in the previously mentioned technical article of J. A. Powell, P. Pirouz, and W. J. Choyke. Generally, the cubic structure, with the zincblende structure, is referred to as β-SiC or 3C—SiC, whereas numerous polytypes of the hexagonal and rhombohedral structures are collectively referred to as α-SiC. To our knowledge, only bulk (i.e., large) crystals of the α polytypes, with sufficient quality for beneficial electronic devices, have been grown to date. The β (or 3C) polytype can only be obtained as (1) small (less than 1 cm2) blocky crystals or thick epitaxial films on small (less than 1 cm2) 3C substrates or as (2) crystal films grown heteroepitaxially on some other substrate. The most commonly available α-SiC polytypes are 4H—SiC and 6H—SiC; these are commercially available as polished wafers that are cut from large crystal boules grown by a sublimation process, described herein. Commercial wafers are presently available up to 100 mm in diameter. But these commercial wafers have a high density of structural imperfections (i.e. dislocation defects) to be described herein more fully. Each of the SiC polytypes has its own specific advantages over the others. For example, (1) 4H—SiC has a significantly higher electron mobility compared to 6H—SiC, (2) 6H—SiC is used as a substrate for the commercial fabrication of GaN blue light-emitting diodes (LEDs), and (3) 3C—SiC has a high electron mobility similar to that of 4H and may function over wider temperature ranges, compared to the α polytypes, but 3C—SiC crystals of sufficient quality and size for beneficial electronic devices have not been available.
The WBG semiconductors, SiC and the III-Nitride compounds, exhibit tetrahedral bonding. In this configuration, each atom of one element is surrounded by (and bonded to) four atoms of the other element type that form the corners of a tetrahedron. Silicon carbide polytypes are formed by the stacking of double layers (also referred to herein as bilayers) of Si and C atoms. Each double layer may be situated in one of three positions known as A, B, and C to be described hereinafter. The sequence of stacking determines the particular polytype; for example, the repeat sequence for 3C is ABCABC . . . (or ACBACB . . . ) the repeat sequence for 4H is ABACABAC . . . . The repeat sequence for 6H is ABCACBABCACB . . . . The number in the polytype designation gives the number of double layers in the repeat sequence and the letter denotes the structure type (cubic, hexagonal, or rhombohedral). The SiC polytype 3C—SiC has a crystal structure that is commonly known as the zinc blende crystal structure, and the SiC polytype 2H—SiC (with a repeat sequence of ABAB . . . ) has a crystal structure commonly known as the wurtzite crystal structure. The stacking direction is designated as the crystal c-axis; it is perpendicular to the basal plane which is the crystal (0001) plane (denoted by the Miller index for the crystal plane). The use of Miller indices to describe various crystal directions and planes is well known to those skilled in the art. Any crystal direction perpendicular to the c-axis is designated as an a-direction. The SiC polytypes are polar in the c-axis directions; in one direction, the crystal face is terminated with silicon (Si) atoms (in this direction, the c-axis is denoted by <0001>, the Miller index for this crystal direction); in the other direction, the crystal face is terminated with carbon (C) atoms (in this direction, the c-axis is denoted by <000 1>). These two faces of the (0001) plane are known as the Si-face and C-face, respectively.
The 3C—SiC (i.e., cubic) polytype has four equivalent stacking directions, and thus there are four equivalent planes, the set of four {111} planes, that are basal planes. Herein, any of the set {111} planes shall be referred to as a (111) plane. Any (111) plane of the cubic structure is equivalent to the (0001) plane of the α polytypes. Herein, the term “cubic crystal” shall also refer to crystals with the zincblende structure and to crystals with the diamond cubic structure. As used herein, “basal plane” shall refer to either the (0001) plane for a α-SiC, or the (111) plane of 3C—SiC or the (111) plane of any crystal with the cubic structure. The term “vicinal (0001) wafer” shall be used herein for wafers whose polished surface (the growth surface) is misoriented less than 10° from the basal plane. The term “mesa” is meant to represent an isolated growth region. The angle of misorientation of a crystal surface from the (0001) plane shall be referred to herein as the tilt angle. The term “homoepitaxial” shall be referred to herein as epitaxial growth, whereby the film and the substrate (wafer) are of the same polytype and material, and the term “heteroepitaxial” shall be referred to herein as epitaxial growth whereby the film is of a different polytype or material than the substrate.
Current theories explaining epitaxial single-crystal growth are well known. Crystal growth can take place by several mechanisms. Two of these are: (1) growth can take place by the lateral growth of existing atomic-scale steps on the surface of a crystal substrate and (2) growth can take place by the formation of two-dimensional atomic-scale nuclei on the surface followed by lateral growth from the steps formed by the nuclei. The lateral growth from steps is sometimes referred to as “step-flow growth.” In the first mechanism, growth proceeds by step-flow from existing steps without the formation of any two-dimensional nuclei (i.e., without 2D nucleation). In the nucleation mechanism, the nucleus must reach a critical size in order to be stable; in other words, a potential energy barrier must be overcome in order for a stable nucleus to be formed. Contamination or defects on the substrate surface can lower the required potential energy barrier at a nucleation site.
Among the structural defects that are observed in commercial 4H— and 6H—SiC wafers are the following: micropipes, dislocations, and low-angle grain boundaries. Dislocations that are observed include: screw dislocations, basal plane dislocations, and threading edge dislocations. As is known to those skilled in the art, the Burgers vector of a screw dislocation (herein also called an axial screw dislocation) is parallel to the line of the dislocation, and the Burgers vector of an edge dislocation is perpendicular to the line of the dislocation.
As described in the previously mentioned technical articles of M. Dudley et al. and U.S. Pat. No. 6,783,592B2, micropipes are a special case of an axial screw dislocation with a large Burgers vector (i.e. more than 2 to 3 times the c-axis repeat distance for a given polytype). They are tubular hollow-core dislocations approximately parallel to the crystal c-axis. This defect is sometimes referred to as a “killer defect” because it effectively shorts out junction devices at low voltage levels. Recently, the density of micropipes has been reduced significantly in commercial SiC wafers, suggesting that micropipes probably will not be a long-range problem. Typical present-day micropipe densities in commercial wafers are in the range 1 to 30 cm−2. The recent reduction in density and the present-day micropipe densities are more fully described in the technical article of A. R. Powell, J. J. Sumakeris, R. T. Leonard, M. F. Brady, St. G. Muller, V. F. Tsvetkov, H. McD. Hobgood, A. A. Burk, M. J. Paisley, R. C. Glass, and C. H. Carter, Jr., “Status of 4H—SiC Substrate and Epitaxial Materials for Commercial Power Applications”, pp. 3-14, Silicon Carbide 2004—Materials, Processing and Devices, Mat. Res. Soc. Symp. Proc. Vol. 815, Editors: M. Dudley, P. Gouma, T. Kimoto, P. G. Neudeck, S. E. Saddow, Materials Research Society, 2004.
Elementary screw dislocations are structurally similar to micropipes, but have a Burgers vector equal to one c-axis repeat distance for a polytype and have a solid (instead of hollow) core. The total density of all kinds of dislocations is about 104 cm−2 in commercial SiC wafers. This includes: axial screw dislocations (103-104 cm2), threading edge dislocations (103-105 cm2), and basal plane dislocations (101-105 cm−2). These different types of dislocations are more fully described in the technical articles of H. Lendenmann, F. Dahlquist, J. P. Bergman, H. Bleichner, and C. Hallen, “High-Power SiC Diodes: Characteristics, Reliability and Relation to Material Defects,” Mat. Sci. Forum, Vols. 389-392, pp. 1259-1264, 2002 (presented at ICSCRM 2001), and S. Ha, P. Mieszkowski, M. Skowronski, and L. B. Rowland, “Dislocation Conversion in 4H Silicon Carbide Epitaxy,” Journal of Crystal Growth, vol. 244, pp. 257-266, 2002.
During SiC crystal growth in the c-axis direction, each axial screw dislocation generates a growing spiral of atomic-scale steps on the growth surface. These steps become preferred deposition sites with the result that growth occurs by step-flow growth as the spiral of steps continually grow outward from the screw dislocation. The growth associated with screw dislocation is further described in the technical article of H. Matsunami, “Technological Breakthroughs in Growth Control of Silicon Carbide for High Power Electronic Devices,” Jpn. J of Appl. Phys., vol. 43, pp. 6835-6847, 2004. In fact, the screw dislocation replicates the crystal stacking information for growth of a specific SiC polytype in the c-axis direction. For basal-plane surfaces with multiple screw dislocations, the spiral steps from each screw dislocation must coalesce perfectly in order avoid the formation of additional defects. This is not the case for pure edge dislocations, and basal plane dislocations. These latter two dislocations cannot provide the stacking information for crystal growth and do not generate spiral growth steps. However, their presence can have a deleterious effect on electronic devices fabricated from the material which is further described in the technical article by R. Singh, “Reliability and Performance Limitations in SiC Power Devices,” Microelectronics and Reliability, vol. 46, pp. 713-730, 2006.
The term “defect free” shall be referred to herein as a single crystal that is free of extended structural defects, such as dislocations that propagate over numerous atoms in at least one direction. The term “defect free” is not meant to describe isolated point defects that involve at most 1 or 2 atoms at an isolated 3D point in the crystal, such as atomic vacancy point defects, interstitial point defects, and impurity point defects.
To date, variations of the physical vapor transport (PVT) (also called the modified Lely process) have been the primary method that has been employed to mass-produce and commercialize electronic quality single-crystal SiC wafers compatible with standard semiconductor wafer fabrication equipment and methods. The modified Lely process is further described in the previously mentioned technical article of A. R. Powell et al, as well as the technical article of K. Semmelroth, N. Schulze, and G. Pensl, “Growth of SiC polytypes by the physical vapour transport technique,” Journal of physics: Condensed Matter, vol. 16, pp. S11597, 2004. In this process, SiC powder is sublimed at approximately 2400° C. and condensed on the surface of a large-area seed crystal (and the subsequent growing boule) which is at a slightly lower surface temperature. Temperature gradients of the order of 25° C./cm are used as the driving force for deposition. This method for growing large reproducible α-SiC crystals relies on the presence of a high density of screw dislocations distributed over the growth surface of the large-area seed crystal to provide growth steps to grow the crystal along the c-axis in order to maintain a uniform polytypic structure. The mechanism of screw-dislocation assisted crystal growth is well known and documented in prior art. In addition to screw dislocations, the α-SiC crystals produced by the seeded sublimation method are also plagued with numerous edge-type dislocations as well as dislocations that lie along the basal plane of the crystal.
The term “Lely seed crystal” is meant to represent a platelet SiC crystal grown by the Lely process and is more fully described in U.S. Pat. No. 2,854,364 of Jan Anthony Lely. Lely seed crystals are typically less than 1 cm2 in area and less than 2 mm in thickness; they also typically have much lower extended defect density than commercial SiC wafers, but the growth process so far has not been suitable for producing the large-diameter wafers needed for commercial device mass production.
A method known as high temperature chemical vapor deposition (HTCVD) may be used for growing large boules of SiC. Instead of a solid source of SiC, the method uses gaseous sources. This method is further described in U.S. Pat. No. 5,704,985 of O. Kordina et al and a technical article of B. Sundqvist, A. Ellison, A. Jonsson, A. Henry, C. Hallin, J. P. Bergman, B. Magnusson, and E. Janzen, “Growth of High Quality p-type 4H—SiC Substrates by HTCVD,” Mat. Sci. Forum, Vols. 473-476, pp. 21-24, 2003 (presented at ECSCRM 2002). The method has an advantage over the PVT method in that composition of the delivered material does not change as a function of time through the growth process. But it still has the problem of using a large-area seed crystal with a high density of extended defects in order to enable commercially-viable growth rates. The result is single crystal boules with a high density of extended defects. Other processes using the approach of HTCVD have been suggested, but to our knowledge, they all have the problem of using large-area seed crystals with a high density of extended defects (i.e. screw dislocations) to enable growth in the c-axis direction.
A two-stage process is known whereby in a first stage of growth, a Lely seed crystal (6-10 mm in diameter), with low-defect density, is grown laterally (in a-directions) to form a larger diameter (25-50 mm) thin crystal by a sublimation process under near isothermal conditions. In a second stage of the two-stage process, sublimation growth in the c-axis direction produces a large boule. This two-stage process is further described in U.S. Pat. No. 5,746,827 of D. L. Barrett et al. However, in a technical article of N. Schulze, D. L. Barrett, and G. Pensl “Near-Equilibrium Growth of Micropipe-Free 6H—SiC Single Crystals by Physical Vapor Transport,” Applied Physics Letters, vol. 72, pp. 1632-1634, 1998 it was reported that near-equilibrium conditions were required to produce micropipe-free crystal growth on low-defect Lely seed crystals. This required a low temperature gradient (less than 5° C./cm) as the driving force for the sublimation growth. This resulted in a maximum growth rate of about 0.3 mm/h for 6H—SiC which is much lower than can be achieved with the standard commercial PVT process for 6H—SiC. It appears that SiC boule growth based on sublimation growth on low-defect seed crystals as taught by U.S. Pat. No. 5,746,827 will require much longer-duration growth runs than current commercial processes. This would render this approach to be more expensive and commercially uncompetitive. A technical article of M. Tuominenm, R. Yakimova, A. Vehanen, and E. Janzen, “Defect Origin and Development in Sublimation Grown SiC Boules,” Materials Science and Engineering B57, pp. 228-233, 1999 also reported poor SiC crystal growth results using Lely seed crystals. In using Lely seed crystals (with a low-density of screw dislocations), several dislocations become dominant resulting in multiple hillocks with valleys in between. Many defects were reported to form in the valleys between the hillocks as more fully described in the technical article of M. Tuominenm et al.
A SiC bulk growth process is known, whereby growth takes place on large-area seed crystals in a {03 38} crystal direction. This process is more fully described in the technical article of K. Nakayama, Y. Miyanagi, H. Shiomi, S. Nishino, T. Kimoto, and H. Matsunami, “The Development of 4H—SiC {03 38} Wafers,” Mat. Sci. Forum, Vols. 389-393, pp. 123-126, 2002 (presented at ICSCRM 2001). This does yield boules that contain small regions with reduced density of extended defects. However, most of the resulting boule has a high density of stacking faults and other extended crystal defects.
A two-stage sublimation process is known for producing bulk SiC crystals with regions of reduced defect density. This process is further described in the technical article of E. N. Mokhov, M. G. Ramm, M. S. Ramm, A. D. Roenkov, Y. A. Vodakov, S. Y. Karpov, Y. A. Makarov, and H. Helava, “Growth of Faceted Free-Spreading SiC Bulk Crystals by Sublimation,” Mat. Sci. Forum, Vols. 433-436, pp. 29-32, 2003, (presented at ECSCRM 2002), as well as in U.S. Pat. No. 6,428,621B1 of Y. A. Vodakov et al. In their process, growth conditions in a first stage produce crystal growth in directions both axial (normal to a planar seed crystal) and lateral relative to the seed crystal. This creates two zones of growth: a central axial core region (due to axial growth) approximately the diameter of the seed crystal with a defect density approximately the same as the seed crystal and a peripheral zone (due to the lateral growth) with a reduced defect density. The first stage of growth is followed by a second stage in which growth is predominately axial (i.e. lateral growth is suppressed). This second stage produces growth in the vertical axial direction with a constant lateral diameter (equal to the larger diameter resulting from the first stage growth). This second stage growth contains (1) a continuation of the core region with a defect density approximately the same as the seed crystal, and (2) a continuation of the reduced density in a peripheral region around the core region with reduced defect density. This process suffers from the same problem described in the previous sections where growth of SiC in the c-axis direction depends on the existence of multiple step sources (i.e. multiple axial screw dislocations). If the coalescence of growth steps from multiple step sources is not perfect (and indeed, this appears to be the case), then defects will be produced by the imperfect coalescence the growth steps in both the core and peripheral regions.
A recently published multi-stage (in different growth directions) process has produced single crystal boules of SiC with a much lower dislocation density and is discussed in the technical article of D. Nakamura, I. Gunjishims, S. Yamaguchi, T. Ito, A. Okamoto, H. Kondo, S. Onda, and K. Takatori, entitled “Ultrahigh-Quality Silicon Carbide Single Crystals,” and published in Nature, vol. 430, pp. 1109-1012, 2004. In each stage of this process, growth starts with a large-area planar seed crystal. By cutting each successive boule in different crystallographic directions (e.g. at each successive stage, a new planar wafer-shaped seed crystal is cut approximately perpendicular to the seed crystal of the previous stage). The final growth stage returns to growth nearly parallel to the c-axis direction. This process results in significant reduction of dislocations in each successive stage except for the last stage that is grown approximately in the c-axis direction. However, in an oral presentation at the European Conference on Silicon Carbide and Related Materials (ECSCRM) Aug. 31-Sep. 4, 2004, at Bologna, Italy, D. Nakumura reported that the growth rate in the c-axis direction in the final growth stage of this new process must be reduced significantly in order to maintain the higher quality of the resulting final boule. This is consistent with the work reported in the previously mentioned technical article of N. Schulze et al and, thus, gives credence to D. Nakumura's oral report. One can conclude that the far fewer remaining screw dislocations must supply spiral steps for step flow growth over much larger areas of the top growth surface compared to conventional SiC boules that are grown with a much higher screw dislocation density. Hence, only a few screw dislocations will dominate the growth surface. Apparently, from the work discussed in the previously mentioned technical articles of N. Schulze et al and M. Tuominenm et al (both groups used low-defect Lely platelets as seed crystals), seed crystals with multiple dislocations (but much fewer than conventional seed crystals used with the modified-Lely process) produce crystals with micropipes and other unwanted defects if growth is carried out at growth rates comparable to the conventional commercial modified-Lely process. It appears that structural defects will continue to be a problem for growth processes that are dependent on the coalescence of steps from multiple step sources (either a high density or a low density of step sources) that are scattered over the surface of a large-area seed crystal. It may be concluded from prior art experience, that the use of large-area seed crystals with multiple step sources (regardless of the density of the step sources) will continue to cause problems (i.e., undesired dislocations) for any commercially-viable growth process for hexagonal or rhombohedral SiC polytypes.
A method for the growth of 3C—SiC crystals is described in the technical article of S. N. Gorin and L. M. Ivanova, “Cubic Silicon Carbide (3C—SiC): Structure and Properties of Single Crystals Grown by Thermal Decomposition of Methyl Trichlorosilane in Hydrogen,” Phys. Stat. Sol. (b), Vol. 202, pp. 221-245, 1997. These individuals grew crystals of 3C—SiC on a heated rod of graphite using chemical vapor deposition. This produced multiple seed crystals each with numerous twin planes (e.g. planar dislocations). As the seed crystals grew, the twin planes propagated throughout the growing crystals and this resulted in crystals with poor morphology and numerous planar dislocations.
A growth of 3C—SiC films on Si is described in the technical article of H. Nagasawa, K. Yagi, T. Kawahara, and N. Hatta, “Reducing Planar Defects in 3C—SiC,” Chemical Vapor Deposition, vol. 12, pp. 502-508, 2006. These individuals grew 3C—SiC films on “undulant” grooved single-crystal (001) Si substrates using chemical vapor deposition. A high density of planar defects was generated at the film/substrate boundary, and the density decreased as the thickness of the films increased. However, the defect density did not decrease to zero, as defects along certain directions could not be eliminated with this process.
In known processes, thin films of low-defect 3C—SiC can be grown heteroepitaxially on small step-free mesas of α-SiC substrates. These processes are further described in U.S. Pat. Nos. 5,915,194; 6,165,874 and 6,488,771B1. However, these processes are not directly suitable for the growth of large bulk crystals because of low growth rates and extended growth that eventually leads to detrimental coalescence with the surrounding hexagonal substrate material.
No commercially-viable method for mass-producing large, reproducible single-crystals of SiC without detrimental quantities of dislocations exists to our knowledge in the prior art. This same predicament (i.e., no method of producing large single-crystals without detrimental quantities of dislocations) also plagues diamond, III-Nitride compounds and alloys, and other crystal materials with a broad variety of useful applications. A serious shortcoming in prior-art processes for the growth of bulk crystals of WPG materials is the use of large-area seed crystals with multiple extended dislocations propagating parallel to the growth direction.
Other non-commercial processes for growing SiC crystals have been reported. A SiC growth process that relates somewhat to the present invention is that described in U.S. Pat. No. 3,721,732. In this process, the objective was to grow multiple whiskers with circular cross section for use as a reinforcing material when mixed with other materials such as glass, plastic, and metals. In this process, a crystal growth process, known as the vapor-liquid-solid (VLS) process and is more fully described in U.S. Pat. No. 3,346,414, is used. In the VLS process, a liquid droplet on a substrate acts as a solvent for material introduced into the vapor contacting the droplet. For the example given in U.S. Pat. No. 3,721,732, the substrate was graphite or a SiC crystal, the droplet material was 20-μm-diameter particles of iron (Fe), or Fe plus Si, and the vapor contained Si and C. At a temperature of about 1200° C., 10-μm-diameter whiskers grew at a rate of 1 mm/h. Whiskers up to 30 mm in length were reported. According to this reference, whiskers, including those grown from droplets on the (0001) SiC Lely crystal surfaces, were mostly of the cubic SiC polytype. Whiskers of a given hexagonal polytype were achieved when the droplet was placed on prismatic (1 100) or pyramidal (11 20) surfaces of a SiC Lely crystal. Both prismatic and pyramidal surfaces are perpendicular (or inclined) to the SiC basal plane (i.e. the (0001) plane). So, for the experiments described in U.S. Pat. No. 3,721,732, the replication of the proper polytype stacking sequence was provided by the stacking sequence template existing on the prismatic and pyramidal planes. This is consistent with growth results wherein commercial SiC wafers are polished “off-axis” in order to provide the polytypic stacking template to the epilayer growth surface as described in the previously mentioned technical article of H. Matsunami. This is also consistent with the nucleation of 3C—SiC on the step-free basal plane surface of 4H— and 6H—SiC substrates and is more fully discussed in the previously mentioned technical article of Neudeck and Powell.
A subsequent U.S. Pat. No. 4,013,503 by Knippenberg and Verspui states that whisker growth is more pronounced on the prismatic and pyramidal SiC surfaces than on the (0001) basal plane. It is now known that, in the absence of defects, growth can only occur on the basal plane by nucleation of cubic SiC as more fully discussed in the previously mentioned technical article of Neudeck and Powell. Hence, the whisker growth methods described by the two patents by Knippenberg and Verspui do not provide or suggest a method for growing large nearly-defect free single crystal boules of non-cubic SiC polytypes and other materials. The teachings of U.S. Pat. Nos. 3,721,732 and 4,013,503 of W. F. Knippenberg et al also do not provide a method for growing a large-diameter single-crystal boule needed for semiconductor mass production.
The growth of high-quality single-crystal whiskers of 2H—SiC that are without dislocations are described in the technical articles of J. A. Powell, “Crystal Growth of 2H Silicon Carbide”, J. Appl. Phys., vol. 40, pp. 4660-4662, 1969, and W. M. Vetter, W. Huang, P. Neudeck, J. A. Powell, and M. Dudley, “Synchrotron White-Beam Topographic Studies of 2H—SiC crystals,” Journal of Crystal Growth, vol. 224, pp. 269-273, 2001. However, the growth process described is not suitable for the growth of large crystals.
A process for making GaN seed crystals is more fully described in the technical article of Philip R. Tavernier and David R. Clarke, “Progress Toward Making Gallium Nitride Seed Crystals Using Hydride Vapor-Phase Epitaxy,” J. Am. Ceram. Soc., vol. 85, no. 1, pp. 49-54, 2002. These individuals describe a technique for producing seed crystals of GaN for subsequent growth into bulk crystals of GaN. In their approach, a GaN film is grown heteroepitaxially on a sapphire substrate by vapor-phase process; then the GaN film is removed from the substrate. The separated GaN film is used as a seed crystal for further GaN growth. The process as reported produces seed crystal with a high density of defects, such as threading dislocations. These defects are incorporated into any bulk crystal grown from the seed crystal.
A process for growing bulk single crystal boules of silicon (Si) by chemical vapor deposition (CVD) for use in the fabrication of solar cell arrays is more fully described in the technical article of A. Franzosi, L. Giarda, and L. Pelosini, “Process of Deposition of Single Crystal Silicon Directly from the Vapour Phase”, PB82-187444, Commission of the European Communities, Energy, EUR-7093-EN, 1981. An important goal in this work was to produce large boules with a hexagonal cross section. The hexagonal shape would allow the fabrication of large arrays of solar cells with improved packing factors, compared to cells with a circular shape. The process consisted of two steps:
(1) Long (40 cm), single-crystal seed crystals, 4.5 mm in diameter, were grown in a <111> direction using a float-zone growth procedure. This float-zone process utilized a growth-from-the-melt process using liquefied silicon.
(2) Two of the long seed crystals were mounted in a large bell jar with electrical contacts on each end. In this manner the seed crystals could be heated to a suitable temperature by passing an electrical current through the seed crystals. The atmosphere in the bell jar was controlled in a manner that produced lateral growth of single crystal Si on the seed crystals. It was found that by suitable choice of gaseous precursors, seed crystal temperature, and other factors, single crystal Si boules, 2 cm in diameter and with a hexagonal cross section, could be grown. This work was reported 25 years ago and, since then, different improved crystal growth processes for Si have superceded this process.
Although some aspects of the process described in the previously mentioned technical article of Franzosi et al. are somewhat related to the current invention, it does not address important problems encountered in the growth of WBG semiconductors. One important problem is that the process of Franzosi et al. does not provide a viable process for producing a long seed crystal for the WBG semiconductor materials since the float-zone crystal growth process (using a melt of the material being grown) is not applicable to the WBG materials. The reason is that most of the WBG materials cannot be melted at any reasonable pressure. Also, their process does not address the problem of maintaining the proper crystal stacking order along the c-axis crystal direction for the hexagonal polytypes of the materials: SiC, AlN, and GaN. Finally, the fairly large diameter (4.5 mm) of their seed crystal (step no. 1) does not provide for a low-defect seed crystal for step no. 2 if applied to WBG materials.
An impurity doping method, known as site competition, is widely used to control the doping of SiC during growth of epitaxial films in the manufacture of semiconductor devices and is more fully described in U.S. Pat. No. 5,463,978. This principle requires the control of the C/Si ratio at the growth site. This principle is difficult to apply during prior-art commercial growth of bulk SiC crystals because of variations in temperature and precursor supply at various locations on the growth front during growth processes.
All commercial SiC wafers, serving as substrates, to date, contain dislocations distributed randomly across the substrate in average densities that are of the order of thousands per square centimeter of wafer area. All of these dislocations present in the wafer propagate in one form or another into the epitaxial layers making up the high field devices. These dislocation defects are difficult to observe (especially in a production environment), and it is nearly impossible to readily predict their locations on any given wafer so that the device being fabricated cannot practically be patterned and/or placed to avoid the vast majority of these defects. This greatly harms the yield, performance, cost, and commercialization of highly beneficial SiC high-field power switching devices. It is highly desirable to reduce or even eliminate the detrimental performance effects of dislocations in SiC crystals and devices.
The present-day commercial c-axis growth methods for growing large reproducible crystals of the non-cubic polytypes of SiC rely on the presence of axial screw dislocations to provide the proper crystal stacking sequence while providing steps for growth of the crystal along the crystal c-axis. If no continuous source of steps is provided (i.e. from an axial screw dislocation), the existing steps on the growth surface (i.e. the basal plane) grow themselves out of existence and growth stops in the c-axis direction, or the 3C (cubic) polytype of SiC nucleates on the surface. The role of growth steps is more fully discussed in the previously mentioned technical article of Neudeck and Powell. The mechanism of axial-screw-dislocation-assisted crystal growth is well known and documented in prior art. No c-axis growth method for mass producing large, reproducible single-crystals of SiC without axial screw dislocations exists in the prior art. This same predicament (i.e., no method of producing large single-crystals without axial screw dislocations) also plagues the III-Nitrides (e.g. GaN and AlN), and other crystal materials with a broad variety of useful applications.
In conventional SiC bulk crystal growth, the diameter of SiC seed crystals used for growth-from-vapor processes is generally planar and of the same order of, or slightly smaller, than the final diameter of the single crystal boule. Thus, the primary growth (i.e. crystal enlargement) direction generally has been approximately in the c-axis direction, which is the stacking direction of Si—C double layers for the non-cubic polytypes. In order to achieve reasonably high growth rates (the order of 0.5 mm/hour or greater) and to promote proper hexagonal stacking sequence throughout the crystal, a fairly high density of screw dislocations (greater than 103 cm−2) has been required. As the screw dislocation density in the seed crystal is reduced, it has been found that crystal growth rate must also be reduced in order to maintain the lower defect density in the growing crystal in a manner more fully described in the previously mentioned technical article of N. Schulze, D. L. Barrett and G. Pensl. Thus, the production rate of crystals with significantly lower screw dislocation density is decreased, and the cost per wafer will be disadvantageously increased using conventional prior art growth techniques.